Method for controlling aluminum titanate ceramic filter properties

ABSTRACT

A method for improving the thermo-mechanical properties of an aluminum-titanate composite, the composite including at least one of strontium-feldspar, mullite, cordierite, or a combination thereof, including:
         combining a glass source and an aluminum-titanate source into a batch composition; and   firing the combined batch composite composition to produce the aluminum-titanate composite.       

     Another method for improving the thermo-mechanical properties of the composite dips a fired composite article into phosphoric acid, and then anneal the dipped composite article. The resulting composites have a thin glass film situated between the ceramic granules of the composite, which can arrest microcrack propagation.

The entire disclosure of any publication or patent document mentionedherein is incorporated by reference.

FIELD

The disclosure relates generally to manufacturing methods forcontrolling aluminum titanate ceramic filter properties, such as for usein combustion exhaust emission control systems.

BACKGROUND

Various manufacturing methods are known for preparing aluminum titanateceramic filters. The present disclosure provides manufacturing methodsfor controlling the physical properties of the resulting aluminumtitanate ceramic filter.

SUMMARY

The disclosure provides methods for controlling the physical properties,such as strength of aluminum titanate (AT) ceramic filters, such as ahoneycomb body, and methods of making and use of the substrates orfilters.

BRIEF DESCRIPTION OF THE DRAWINGS

In embodiments of the disclosure:

FIGS. 1A to 1F summarize the comparative toughening of microcrackedceramics having thin intergranular glass.

FIGS. 2A to 2D show SEM images of a polished cross section of anexemplary ceramic material having lithium sourced intergranular glass.

FIGS. 3A to 3D show SEM images at two different magnifications ofpolished cross sections of ceramic materials that are free of lithium(FIGS. 3A and 3B) and have lithium sourced intergranular glass (FIGS. 3Cand 3D).

FIG. 4 shows pore size distributions for materials A (control), B, and Chaving, respectively, 0%, 0.5%, and 1 wt % Li-source addition.

FIGS. 5A and 5B show thermal expansion curves for fired ceramicmaterials without and with a lithium-source in the batch.

FIGS. 6A and 6B show examples of fired ceramic elastic modulushysteresis upon thermal cycling.

FIGS. 7A to 7C show composited SEM micrographs of three fired ceramicsamples at three different magnification levels (left to right) thatwere prepared with increasing levels of aluminum-source (aluminumphosphate) additive.

FIGS. 8A and 8B show differential scanning calorimetry results forselected fired ceramic materials.

FIGS. 9A and 9B show aspects of the pore size distribution properties offired ceramic samples made from batches having different levels ofphosphorous-source addition.

FIG. 10 shows thermal expansion curves for fired ceramic samples frombatches A, D, E, and F having different levels of phosphate-sourceaddition.

FIGS. 11A to 11C show pore size distributions for fired ceramic samplesbefore and after H₃PO₄ dipping, then annealing.

FIGS. 12A to 12F show SEM images of polished cross sections (FIGS. 12Ato 12C) and surfaces (FIGS. 12D to 12F) of as-fired ceramic parts andthese same ceramic parts after dipping in phosphoric acid and thenannealing.

FIGS. 13A, 13B, and 13C, respectively, show CTE results for originalfired ceramic parts and phosphoric acid dipped fired ceramic parts foreach of samples A, G, and H.

FIGS. 14A, 14B, and 14C show the hysteresis of the elastic modulusduring heating and cooling of fired ceramic parts and phosphoric aciddipped fired ceramic parts for samples A, H, and G.

FIG. 15 shows pore size distribution results for aluminumtitanate-strontium feldspar composite samples having varying glasscontent.

FIGS. 16A and 16B show as-measured MOR (FIG. 16A) andporosity-normalized MOR (16B) of selected aluminum titanate-strontiumfeldspar composites having different glass levels.

FIG. 17 shows pore size distribution properties ofalumina-cordierite-mullite (MgO) composites having boron oxide additionin the batch and compared to a boron-oxide free control.

DETAILED DESCRIPTION

Various embodiments of the disclosure will be described in detail withreference to drawings, if any. Reference to various embodiments does notlimit the scope of the invention, which is limited only by the scope ofthe attached claims. Additionally, any examples set forth in thisspecification are not limiting and merely set forth some of the manypossible embodiments of the claimed invention.

DEFINITIONS

“Porosity,” and like terms generally refer to the void spaces, orsynonymously, pores within the walls of the honeycomb material. The voidspace in a honeycomb occupied by the macroscopic channels is excluded.“Porosity,” and like terms generally refer to the total void space in ahoneycomb material that can be attributed to the presence of pores andexcludes the void space in a honeycomb material attributable to thepresence of macroscopic channels or vias of the honeycomb, or the ratioof the pore volume to the total volume of a pulverized solid material,and may be expressed as percent porosity (% P). Porosity, and likeaspects of the ceramic bodies, are mentioned in commonly owned andassigned U.S. Pat. No. 6,864,198. Parameters such as d10, d50 and d90relate to the pore size distribution. The quantity d50 is the medianpore size (MPS) based upon pore volume, and is measured in micrometers;thus, d50 is the pore diameter at which 50% of the open porosity of theceramic has been intruded by mercury (mercury porosimetry). The quantityd90 is the pore diameter at which 90% of the pore volume is comprised ofpores whose diameters are smaller than the value of d90; thus, d90 isequal to the pore diameter at which 10% by volume of the open porosityof the ceramic has been intruded by mercury. The quantity d10 is thepore diameter at which 10% of the pore volume is comprised of poreswhose diameters are smaller than the value of d10; thus, d10 is equal tothe pore diameter at which 90% by volume of the open porosity of theceramic has been intruded by mercury. The values of d10 and d90 are alsoin units of micrometers. The quantity (d50−d10/d50) describes the widthof the distribution of pore sizes finer than the median pore size, d50.

“Super additive,” “super addition,” and like terms generally refer toadding additional ingredients or materials to a batch composition orlike formulation in excess of, or in addition to, a 100 wt % baseinorganics formulation. A base formulation totaling 100 wt % can be, forexample, a combination of ceramic forming ingredients in an amount from20 to 70 weight percent and an inorganic filler material in an amountfrom 80 to 30 weight percent, and the super additives can be, forexample, one or more pore formers, a glass former source material, or acombination thereof, with or without other super additives, and can bepresent or added to the batch in, for example, from about 50 to about300 wt % in addition to the base formulation 100 wt %.

“Include,” “includes,” or like terms means encompassing but not limitedto, that is, inclusive and not exclusive.

“About” modifying, for example, the quantity of an ingredient in acomposition, concentrations, volumes, process temperature, process time,yields, flow rates, pressures, and like values, and ranges thereof,employed in describing the embodiments of the disclosure, refers tovariation in the numerical quantity that can occur, for example: throughtypical measuring and handling procedures used for making compositions,concentrates, or use formulations; through inadvertent error in theseprocedures; through differences in the manufacture, source, or purity ofstarting materials or ingredients used to carry out the methods; andlike considerations. The term “about” also encompasses amounts thatdiffer due to aging of a composition or formulation with a particularinitial concentration or mixture, and amounts that differ due to mixingor processing a composition or formulation with a particular initialconcentration or mixture. The claims appended hereto include equivalentsof these “about” quantities.

“Consisting essentially of” in embodiments refers, for example, to acatalytic honeycomb filter article having relatively high disclosedporosity and increased strength, to a method of making a catalyticfilter article and precursors thereto, devices incorporating thecatalytic filter article, and can include the components or steps listedin the claim, plus other components or steps that do not materiallyaffect the basic and novel properties of the compositions, articles,apparatus, or methods of making and use of the disclosure, such asparticular reactants, particular additives or ingredients, a particularagents, a particular surface modifier or condition, or like structure,material, or process variable selected. Items that may materially affectthe basic properties of the components or steps of the disclosure orthat may impart undesirable characteristics to the present disclosureinclude, for example, an article having significantly reduced porosity,and little or no improvement in strength or microcrack propagationcontrol of the article, and that are beyond the values, includingintermediate values and ranges, defined and specified herein.

The indefinite article “a” or “an” and its corresponding definitearticle “the” as used herein means at least one, or one or more, unlessspecified otherwise. Abbreviations, which are well known to one ofordinary skill in the art, may be used (e.g., “h” or “hr” for hour orhours, “g” or “gm” for gram(s), “mL” for milliliters, and “rt” for roomtemperature, “nm” for nanometers, and like abbreviations).

Specific and preferred values disclosed for components, ingredients,additives, and like aspects, and ranges thereof, are for illustrationonly; they do not exclude other defined values or other values withindefined ranges. The compositions, apparatus, and methods of thedisclosure can include any value or any combination of the values,specific values, more specific values, and preferred values describedherein. Cordierite and aluminum titanate (AT)-based honeycomb substrateshave been used for a variety of applications such as catalytic convertersubstrates and filters for diesel particulate emission. In response tothe increasingly restrictive emission standards for light and heavy dutyvehicles, the filter materials need to be highly porous to allow gasflow through the walls without compromising the engine power; show highfiltration efficiency for emitted particles and not suffer majorpressure drop. The filters also have to withstand the corrodingenvironment of the exhaust gas and be able to bear severe thermal shockduring rapid heating and cooling. High material green strength and highintermediate firing strength are desired to achieve high precision inextruded shape and high firing selects and also allows implementation offaster firing cycles with associated energy savings. High mechanicalrobustness is also needed to withstand the mechanical constraints duringfilter handling and canning.

Cordierite and aluminum titanate-based materials have low thermalexpansion and are therefore suited for applications having high thermalshock resistance. Both materials show strong anisotropy in their thermalexpansion with different crystallographic directions exhibiting positiveand negative expansion. Due to the anisotropy in thermal expansion,mismatch strains build up between grains with different crystallographicorientation and can lead to microcracking. Polycrystalline cordieriteand aluminum titanate ceramics can undergo extensive microcrackingduring thermal cycling.

Microcracks open during cooling and close, sometimes even heal, duringheating. This creates a hysteresis response to thermal cycling withdifferences between the heating and cooling curve that can be attributedto the presence of microcracks. As a result of the microcracking, theoverall thermal expansion of the ceramic piece is lowered compared tothe crystallographic average CTE.

On first impression, lowering of CTE through microcracking seemsbeneficial; the thermal shock resistance of the material, which isproportional to the material's strength and inversely proportional toits elastic modulus and thermal expansion, is expected to be improved.However, material strength also decreases with increasing microcrackdensity. Microcrack densities in cordierite remain rather low, due tothe small difference in crystallographic thermal expansion, and largegrain (domain) sizes that are needed to reach the microcracking stressthreshold. As a result of the larger anisotropy in crystallographicexpansion, microcrack densities in aluminum titanate-based materials canbe very high and strongly limit the ceramic's strength. Due to the veryhigh anisotropy and high absolute values in thermal expansion, themicrocrack density is always high in aluminum titanate-based materials,but differs from material to material depending on, for example,aluminum titanate grain size, local mis-orientation of adjacent grains,medium range orientation arrangement of the grains (domains), and thenature and distribution of second phases. Since thermal expansion,strength, and porosity of aluminum titanate-based materials are highlycoupled through the microcrack density, it remains a challenge to makehighly porous aluminum titanate-based honeycomb ceramic articles thatcombine low thermal expansion, high porosity, low Young modulus, highstrength, and are attractive for high-performance diesel particulatefilter applications. Tightening of regulations toward higher filtrationefficiency, further reduced CO₂ emission, and less fuel consumptionmotivate automakers to use filters with the lowest possible pressuredrop at the highest possible filtration efficiency, both at improvedthermal shock resistance and extended lifetime. Such filterspecifications can impose, for example, higher porosity, thinnerhoneycomb walls, or both for filters, which require significantimprovements in material strength at a given pressure drop andfiltration efficiency.

Several approaches for material strength improvement have beendemonstrated in the past for aluminum titanate-based ceramics, forexample:

1. The aluminum titanate grain size in the material can be reduced,leading to a lower microcrack density in the material and an increase inmaterial strength. The concept was implemented in aluminumtitanate-based materials made from fine inorganic alumina fiber with aresulting aluminum titanate grain size of about 3 to 7 micrometers onmicrostructures and fibrous AT materials (see commonly owned andassigned copending application U.S. Ser. No. 61/067,615, filed Feb. 29,2008, entitled “Acicular Porous Ceramic Article and ManufactureThereof”).

2. Improved thermo-mechanical properties were also implemented byTepesch et al., by choice of a narrow raw material (alumina) particlesize distribution, see commonly owned and assigned copending applicationUS20100222200A1 entitled “Aluminum Titanate-Containing Ceramic-FormingBatch Materials And Methods Using The Same”.

3. Aluminum titanate in the material can be textured so that the overallmicrocrack density is reduced. Such texturing was realized in the abovementioned copending application U.S. Ser. No. 61/067,615, wheretemplated growth of aluminum titanate on inorganic alumina fiber led toa preferential alignment of the negative expansion c-axis of growingaluminum titanate grains in the fiber axis and, since fibers alignedupon extrusion in the extrusion axis, also resulted in the texture ofthe aluminum-titanate in the honeycomb axis.

4. Domain size can be kept small so that the local stresses remain smalland anisotropic, and the overall microcrack density is kept low.Extremely small domain size and high strength have been demonstrated inadvanced aluminum-titanate-type materials with domain sizes of about 20micrometers and less. A domain refers to an entity of neighboring grainsthat do not exceed a mis-orientation of 15 degrees in their negativeexpansion c-axis (see the abovementioned U.S. Ser. No. 61/067,615).

In embodiments, the disclosure provides a method for improving thethermo-mechanical properties of an aluminum-titanate composite, thecomposite including at least one of strontium-feldspar, mullite,cordierite, or a combination thereof, comprising:

combining a glass source and an aluminum-titanate source into a batchcomposition; and

firing the combined batch composition to produce the aluminum-titanatecomposite,

wherein the glass source provides a glass film, a glass-ceramic film, orboth, between the ceramic granules of the composite.

The glass source, such as a sintering aid, can be, for example, at leastone of a lithium oxide, a boron oxide, a silica, a phosphorus oxide, aprecursor or source of any of the foregoing, and like materials, ormixtures or combinations thereof. The aluminum-titanate source batchcomposition can be, for example, at least one of an aluminum source anda titania source. Sources of alumina can include, for example, powdersthat when heated to a sufficiently high temperature in the absence ofother raw materials, will yield substantially pure aluminum oxide.Examples of such alumina sources include alpha-alumina, a transitionalumina such as gamma-alumina or rho-alumina, hydrated alumina,gibbsite, corundum (Al₂O₃), boehmite (AlO(OH)), pseudoboehmite, aluminumhydroxide (Al(OH)₃), aluminum oxyhydroxide, and mixtures thereof. Inembodiments, the at least one alumina source can be, for example, atleast 40 wt %, at least 45 wt %, or at least 50 wt % of the inorganicmaterials, such as, for example, 49 wt % of the inorganic materials. Inembodiments, the at least one alumina source can be selected so that themedian particle diameter of the at least one alumina source is from 1 to45 microns, for example, from 2 to 25 microns.

Sources of titania can include, for example, rutile, anatase, strontiumtitanate, titanium silicates, amorphous titania, and mixtures thereof.In embodiments, the at least one titania source can comprise at least 20wt % of the inorganic materials, for example, at least 25 wt % or atleast 30 wt % of the inorganic materials. In embodiments, the inorganicmaterials can further comprise at least one additional material. Inembodiments, the at least one additional material can be chosen fromsilica, oxides (e.g., lanthanum oxide), carbonates (e.g., calciumcarbonate and strontium carbonate), nitrates, and hydroxides. Inembodiments, the at least one additional material can be silica, whichcan further comprise at least 5 wt % of the inorganic materials, forexample, at least 8 wt % or at least 10 wt % of the inorganic materials.

Sources of silica include, for example, non-crystalline silica, such asfused silica or sol-gel silica, silicone resin, low-aluminasubstantially alkali-free zeolite, diatomaceous silica, kaolin, andcrystalline silica, such as quartz or cristobalite. Additionally, thesources of silica may include silica-forming sources that comprise acompound that forms free silica when heated, for example, silicic acidor a silicon organometallic compound. In embodiments, mullite or otheraluminum silicates, or more complex silicates, can also be a mixedalumina-silica source. Thus, the alumina source can contain otherconstituents of the final composite.

The combining of the glass source can be accomplished by, for example,at least one of:

adding a combination of silica and an alkali oxide source, such aslithium oxide, sodium oxide, potassium oxide, or their precursors, suchas hydroxides, halides, and like sources, to the batch in an amount of0.3 to 5 wt %, preferably 0.3 to 3 wt %, or more preferably 0.4 to 1 wt%;

adding a phosphorous oxide source to the batch in an amount of 0.1 to 3wt %, and preferably 0.1 to 1%;

adding a boron oxide source to the batch in an amount of 0.1 to 2 wt %,and preferably 0.1 to 1%;

or a combination thereof, the weight % can be based on a superadditionrelative to the total weight of the batch composition.

The glass source can produce a glass or glass-ceramic phase during anearly stage of firing the batch, and the glass or glass-ceramic phasetransforms during a later stage of firing the batch, the transformationcomprising at least one of:

a partial glass or glass-ceramic phase evaporation;

a glass or glass-ceramic phase crystallization;

a glass or glass-ceramic phase incorporation into the composite phase;

or a combination thereof.

As examples, Li sources can partially evaporate, B sources can partiallyevaporate and leave the residual glass having a different composition, Psources can redistribute within the composite, and Si—La—Ti sources canpartially crystallize. In embodiments, the combining and firing canimprove the composite strength (porosity-normalized) by, for example, 5to 25% at a CTE less than 10×10⁻⁷K⁻¹ (preferably 3×10⁻⁷K⁻¹) at 50%porosity compared to the composite prepared in the absence of the glasssource. The results are believed to be applicable to other porosities.

In embodiments, the firing can be accomplished, for example, at 1390° C.to 1410° C., for 10 to 20 hours.

In embodiments, at the same top-temperature hold-time:

the firing temperature can be, for example, reduced by at least 25° C.;

the firing time can be, for example, reduced by at least 10%;

or a combination of reduced firing temperature and reduced firing time,to provide the same useable product compared to the method practicedfree of the glass source.

In embodiments, the combining and firing the batch including the glasssource is believed to accelerate aluminum-titanate phase formationduring an initial stage of firing the batch, at lower temperature, inless time, or both, compared to a batch without the included glasssource. The combining and firing the batch including the glass sourcecan provide the composite having 3 to 5 micrometer smaller medianceramic grain size. As a result of the smaller grain size the compositecan have a higher strength compared to a ceramic prepared from a batchwithout the glass source.

In embodiments, the disclosure provides a method for improving thethermo-mechanical properties of an aluminum titanate-based composite,comprising:

dipping a fully fired aluminum titanate-based composite into an aqueousphosphorous acid solution of 0.5 to 10 wt %; and

annealing the dipped composite, to provide phosphorous incorporationinto the resulting composite of 0.5 to 2 wt %, the weight % being basedon a superaddition relative to the total weight of the un-dippedcomposite.

In embodiments, the disclosure provides a method for toughening amicrocracked aluminum titanate ceramic, comprising:

creating an intergranular glass film within the ceramic, the film havinga thickness of from about 20 nm to 500 nm, and the film interacting withthe microcracks and limiting uncontrolled growth of the microcracks.

In embodiments, the ceramic can be toughened from about 5 to about 25%,as demonstrated by, for example, the increase in the modulus of rupturemeasured by 4-point bending relative to a ceramic prepared without theintergranular glass film.

In embodiments, the disclosure provides a method comprising dipping afired composite article into an aqueous phosphorous acid solution of 0.5to 10 wt %, and annealing the dipped article, to provide the compositehaving phosphorous incorporation after firing of 0.5 to 2 wt % based onthe weight of the fired composite.

In embodiments, the disclosure provides a method for improving thethermo-mechanical properties of an aluminum-titanate composite, thecomposite including at least one of strontium-feldspar, mullite,cordierite, or a combination thereof, that comprises, consistsessentially of, or consists of:

combining a glass source and an aluminum-titanate source into a batchcomposition; and

firing the combined batch composite composition to produce thealuminum-titanate composite,

wherein the glass source provides a glass film, a glass-ceramic film, orboth, between the ceramic granules of the composite.

In embodiments, the disclosure provides a method for improving thethermo-mechanical properties of an aluminum titanate-based composite,that comprises, consists essentially of:

dipping a fully fired aluminum titanate-based composite into an aqueousphosphorous acid solution of 0.5 to 10 wt %; and

annealing the dipped composite, to provide phosphorous incorporationinto the resulting composite of 0.5 to 2 wt %, the weight % being basedon a superaddition relative to the total weight of the un-dippedcomposite.

In embodiments, the disclosed compositions and articles thereof, and themethod of making and use provide one or more advantageous features oraspects, for example, as discussed below. Features or aspects recited inany of the claims are generally applicable to all facets of theinvention. Any recited single or multiple feature or aspect in any oneclaim can be combined or permuted with any other recited feature oraspect in any other claim or claims.

In embodiments, the disclosure provides methods of making ceramicarticles, which methods are advantaged by using porous alumina-titanatebased ceramic articles for filter and substrate applications, and moreparticularly to an improved batch composition and process for makingceramic articles employing suitable glass-forming agents in the batch.The glass-forming agents form a glassy phase that accelerates thealuminum titanate formation during an initial firing stage of the batch.At later firing stages, the glass-forming agents can, for example,partially evaporate, crystallize, incorporate into the final productphase, or a combination thereof, to leave a fired ceramic article with athin intergranular glass film and overall improved thermo-mechanicalproperties. The resulting ceramic materials have, for example, improvedstrength, low thermal expansion, and high porosity. The attributes ofthe products can provide higher thermal shock resistance throughimproved fracture toughness (e.g., low thermal expansion, high strength,and low elastic modulus) compared to existing products and knownmethods. The attributes also include an improved pore size distributionhaving no, or a significant decrease of, small size pores. Thesematerial properties can provide, for example, improved filtercharacteristics, lower pressure drop, higher filtration efficiency, andhigher thermal shock resistance.

A disadvantage of a known batch composition without particulate glassformer addition or with a sintering additive, such as lanthanum oxide,is its tendency to form very large, localized glass pockets and whichpockets do not extensively wet the grain boundaries and interfaces inthe composite.

In embodiments, the disclosure provides methods for tougheningmicrocracked aluminum titanate-based ceramics comprising incorporating athin intergranular glass film that interacts with the microcracks andlimits uncontrolled microcrack growth. The method improves the strengthand toughness of the ceramic material and enhances the diesel particlefilter (DPF) thermal shock resistance. The filters can have a wideroperating temperature range.

In embodiments, the present disclosure provides a method for improvingthe thermo-mechanical properties of aluminum titanate-based composites(AT) by, for example, including glass forming aids, that is, anengineered glass phase which can achieve improved material strength atlow thermal expansion and high porosity. Current DuratrapAT® materialsare made of an aluminum titanate-strontium feldspar composite withadditional minor phases, including a glass. The glass in DuratrapAT® isconfined to rare, large size pockets and does not penetrate and wet thegrain boundaries and interfaces. In known AT-based materials,microcracks propagate intragranularly through the aluminum titanate andfeldspar grains, with no apparent preference for either and undergo onlylimited crack interaction with the planar defects within the feldspargrains. No particular or special interaction of the cracks is observedwith the glass pockets in DuratrapAT®.

Although not desired to be limited by theory, it is believed that athin, intergranular glass film (having lower strength than the feldspar)in aluminum titanate-strontium feldspar composites can offer a preferredcrack path for propagating microcracks along the wetted grain boundariesand interfaces, so that the intragranular microcracks in the aluminumtitanate grains propagate through the aluminum titanate grain and thenpreferentially through the intergranular glass and no longer through thefeldspar grains. Interaction between the intergranular glass and thepropagating cracks produce crack deflection, interface debonding, crackbifurcation, etc., thus dissipating energy, so that the crackpropagation is slowed down and eventually stopped. Thus theintergranular glass that surrounds the aluminum titanate grains canlimit or decrease the crack propagation from one aluminum titanate grainto the next. Such confinement or containment of microcracking to thealuminum titanate grain and its surrounding glass film providessignificant enhancement in toughness. The energy barrier for crackgrowth beyond the aluminum titanate grain and its surroundingintergranular glass film is increased and thus long-range crackpropagation and growth of microcracks into detrimental macroscopiccracks is retarded. The result is a higher strength material with highfracture toughness. In embodiments, a composite with a very regularphase distribution and lack of percolation of the aluminum titanate ispreferred and shows the strongest effects. Preferred attributes of theglass are low liquid phase formation temperature, suitable viscosity,and high wettability together with a strength of the resulting glassthat is lower than that of the feldspar. The glass phase should wet thealuminum titanate grains well even at low levels.

Various intergranular glasses can be selected. Their level or amount(wt. %) should remain relatively small, for example, from about 0.3weight % to about 3 weight %, so that the CTE of the material is notsignificantly altered. The glass should be compatible with the aluminumtitanate and feldspar. It is desirable that excess of the glass formingingredient, for example, either evaporates (example of boron oxideduring firing at temperatures >1400° C.) or incorporates into thefeldspar (lithium oxide or phosphorous oxide), or both, so that theamount of intergranular glass is kept very low and a very thinintergranular glass film is obtained.

The glass source (i.e., the glass forming aid) can further introduce anearly onset and also an earlier completion of the aluminum titanateformation in the reactive firing, so that a several micrometer smallermedian aluminum titanate grain size can be obtained by this approach. Inexamples of particular aluminum titanate-based composites where duringthe reactive firing full formation of the aluminum titanate productphase is reached at 1390° C., a rise in the firing temperature by 25° C.produces a doubling in aluminum titanate grain size. While the preciseyield and grain growth evolution can depend on the raw material type,particle size, and on the firing schedule, the general trend of graingrowth during higher temperature sintering is a known phenomenon.Additionally, in the presence of glass formers, shorter or lower toptemperature firing cycles are particularly useful.

In embodiments, the disclosure provides methods of making ceramicarticles, which methods are advantaged by using glass forming sinteringaids, such as lithium oxide, boron oxide, silica, titanio-aluminosilicate, and phosphorus oxide. The examples of lithium oxide and boronoxides have been chosen, in embodiments, as representative examples ofthe class of glass forming sintering-aids, such as alkali-containingglasses, magnesio-alumino-silicates, and like low melting silicateglasses.

In embodiments, the present disclosure provides examples of aluminumtitanate-feldspar composites that contain an intergranular phaseprepared by, for example, combining excess silica with lithium oxide;adding phosphorous oxide with the green batch; adding boron oxide withthe green batch; adding alumino-titanio-silicate with the green batch;dipping the fired ceramic into phosphorous acid, or combinationsthereof.

For additional definitions, descriptions, and methods of siliceousformulations, silica materials and related metal oxide materials, seefor example, R. K. Iler, The Chemistry of Silica, Wiley-Interscience,1979.

In embodiments, the present disclosure provides a number of advantages,including for example:

1. Presence of thin intergranular glass films induce preferred crackpropagation in aluminum titanate-feldspar composites and retard thelong-range crack propagation, to provide improved strength/fracturetoughness to the material with otherwise unaltered physical properties.Improved material strength and toughness (compared to, for example,Corning's current Duratrap AT® material and its porosity-engineeredderivatives) are desired, for example, to attain improved thin wall andhigh porosity filters for integrated SCR applications.

2. Glass forming additives induce an early onset of the aluminumtitanate formation and allow completion of the overall reaction at lowertemperature, shorter time, or both, so that the resulting overallaluminum titanate grain size is smaller and the material strengthimproved.

Referring to the figures, FIG. 1 illustrates the general tougheningprinciple discussed above and achieved by thin intergranular glass filmsin microcracked ceramics (FIGS. 1E to 1F) compared to materials withoutany intergranular glass film (FIGS. 1A to 1B) or thick intergranularglass film (FIGS. 1C to 1D). FIGS. 1A, 1C, and 1E show the schematiccomposite material with its aluminum titanate grains (open or unshaded)and feldspar grains (cross-hatched) prior to application ofthermo-mechanical stress with an existing small microcrack(lllllll-line), FIGS. 1B, 1D, and 1F sketch the material afterapplication of thermo-mechanical stress and propagation of themicrocrack.

FIGS. 1A and 1B schematically show an aluminum titanate (opengrains)-feldspar (cross-hatched grains) composite having a glass pocket(black) before (1A) and after (1B) crack propagation, respectively. Aninitial microcrack (lllllllll-line) in FIG. 1A propagates under load ormechanical stress or thermo-mechanical stress through the aluminumtitanate and feldspar grains shown in FIG. 1B. FIGS. 1A and 1B,respectively, schematically show a two-phase microcracked ceramicwithout any intergranular glass films having an initial small microcrackand propagated microcrack after having suffered thermo-mechanicalstress. The microcrack has propagated through aluminum titanate andfeldspar grains and extended over a large relative area.

FIGS. 1C and 1D schematically show an aluminum titanate (opengrains)-feldspar (cross-hatched grains) composite having thickintergranular glass films (thick borders or grain boundaries) before(1C) and after (1D) crack propagation, respectively. The initialmicrocrack (lllllllll-line) in FIG. 1C propagates under load ormechanical stress or thermo-mechanical stress directly on a short paththrough the thick glass film. The thick film of glass in the grainboundaries behaves like a bulk glass, fractures easily, and allows easycrack propagation, so that the crack propagates through the entire“piece.” The microcrack can propagate indiscriminately through the glassphase and some aluminum titanate grains and extend over a very largerelative area. The microcrack does not encounter substantial dissipationof the strain energy; thus the propagation of the microcrack can be veryextensive and can lead to failure of the piece if the microcrack crossesa major fraction of the piece.

FIGS. 1E and 1F schematically show an aluminum titanate (opengrains)-feldspar (cross-hatched grains) composite having thinintergranular glass films (thin line along the grain boundaries) before(1E) and after (1F) crack propagation, respectively. The initialmicrocrack (lllllllll-line) in FIG. 1E propagates under load, mechanicalstress, or thermo-mechanical stress through the thin glass film,interacting on its long path with the adjacent grains and other glassfilm. For example, crack deflection, crack branching, interfacedebonding, and like phenomena, can be active so that a long crack pathor extended crack pattern is obtained, where considerable “crackingenergy” can be dissipated so that the propagating crack can beterminated.

The thin intergranular glass film does not have the properties of anextended bulk glass. Instead, the thin intergranular glass film can bein a strained and chemically, structurally, or both, altered state bythe adjacent grains so that it possesses very different properties,which can allow for a stronger interaction of the propagating crack withvarious glass-grain interfaces and with the glass film itself. Assumingthat fracture energies for the aluminum titanate grain, the aluminumtitanate interface, and the thin glass film are in a same range ofmagnitude, then the propagating microcrack encounters“energy-dissipating” interactions on its path, for example, by eithercreating further microcracking of aluminum titanate grains or debondingaluminum titanate-glass interfaces or splitting into several smallermicrocracks that propagate along several intergranular films or crackdeflection when encountering a high energy interface. FIG. 1F shows thatthe microcrack can be restricted to a long path within the same volumeof the ceramic. Significantly more energy can be dissipated with suchinteraction of the crack with its surroundings compared to a single orstraight crack so that the crack finally stops and does not lead tomaterial failure. Although not bound by theory, FIGS. 1E and 1Fillustrate what is believed to be the active toughening mechanismachieved by the incorporation of thin intergranular glass film into thedisclosed ceramic articles and methods of making.

EXAMPLES

The following examples serve to more fully describe the manner of usingthe above-described disclosure, as well as to set forth the best modescontemplated for carrying out various aspects of the disclosure. It isunderstood that these examples do not limit the scope of thisdisclosure, but rather are presented for illustrative purposes. Theworking examples further describe how to prepare the porous articles ofthe disclosure.

Preparation of a Green Body.

A green body can be prepared according to U.S. Pat. No. 5,332,703,entitled “Batch Compositions for Cordierite Ceramics,” and U.S. Pat. No.6,221,308, entitled “Method of Making Fired Bodies,” both assigned toCorning, Inc., and as modified according to the present disclosure.Various glass forming additives, ceramic compositions, and the resultingmicrostructures and material properties, are described in the followingexamples and are organized by additive and additive addition, in theorder of: Li₂O, phosphate, dipping in phosphoric acid, addition ofexcess silica, and lastly, the addition of B₂O₃.

Example 1 Addition of Lithium Oxide Glass-Forming Additive to the Batch

Separately batched compositions A, B, C having 0, 0.5 and 1% wt Li₂Osuper-addition listed in Table 1 were ram-extruded into 1″ diameterhoneycomb with a 300/14 cell geometry. The inorganic raw materials, poreformer, and binder were pre-mixed dry. The dry mixed ingredients werecombined in a pan and mulled under addition of batch water until asuitable paste texture was reached. Lithium was added to the batch asdissolved lithium acetate in the batch water. Batch composition AA isanother reference batch having a different pore former level asindicated in Table 1.

TABLE 1 Batch compositions with Li-addition. Batch In wt. % A AA B CINORGANICS Silica 10.19 10.19 10.19 10.19 Strontium Carbonate 8 8 8 8Calcium Carbonate 1.38 1.38 1.38 1.38 Titanium Dioxide 29.95 29.95 29.9529.95 Hydrated Alumina 3.71 3.71 3.71 3.71 Lanthanum Oxide 0.2 0.2 0.20.2 Alumina 45.57 45.57 45.57 45.57 Solid inorganics totals 100 100 100100 PORE FORMERS Graphite 10 10 10 Potato Starch 8 8 8 Corn Starch 15SOLID BINDERS/ORGANICS Methylcellulose 4.5 4.5 4.5 4.5 OTHER LIQUIDADDITIONS Lithium Acetate in 20 mL water 0 0 0.2 0.39 Emulsion T 16 1616 16

The resulting paste was then extruded on a ram extruder into a honeycombshape with a honeycomb die and shim of appropriate size. For 1″ ramextruded parts, the die geometry was (300/14) with 300 cells per squareinch and an extruded wall thickness of 14 mil.

The extruded green honeycomb parts were dried in a microwave oven atmedium power for 5 minutes and then further dried in a drying oven for24 h at 85° C. The parts were then fired in air in a CM furnace(cmfurnaces.com) using a ramp rate of 120° C./h, top temperature 1390,1400, or 1410° C., and a hold time of 15 hr. All materials wereextruded, dried, and fired without any crack formation failure problems.

Microstructure of Fully Fired Materials Containing Li Glass FormingAdditive.

The fully fired materials B and C having 0.5 and 1% Li show typicalfeldspar, aluminum titanate, and alumina distribution, but aredistinguished from Li-free A-material by the presence of very regularlydistributed small size glass pockets with extended grain boundary glassfilm. The reference material A contains occasional glass pockets thatare large and not well distributed. The microcracks appear to runthrough the aluminum titanate phase with Li-glass pockets present. Inthe reference Li-free material A, cracks easily propagate through thefeldspar phase. In the Li-containing sample, the microcracks aresmaller, are less frequently observed in feldspar areas and seem to bestopped at the feldspar interfaces and do not propagate through theentire material.

Electron back-scattered diffraction (EBSD) with a domain size analysis(mis-orientation greater than 15 degrees as distinction criterion fordifferent domains) and grain size analysis (mis-orientation greater than1 degree as distinction criterion for different grains) showed that theaverage aluminum titanate grain size is 10 micrometers and the averagedomain size about 40 micrometers for the material with 1% Li. Noparticular texturing was observed in the pole figure of aluminumtitanate. The microstructural characteristics of aluminum titanate inmaterials derived from Li-containing batches are very similar to thereference material. The aluminum titanate grain size is not largerdespite the glass flux.

FIGS. 2A to 2D show SEM images (electron backscattered diffraction) of apolished cross section of an exemplary ceramic material having lithiumsourced thin intergranular glass; specifically the above experimentalC-material containing 1% Li.

FIG. 2A is a band contrast image showing the typical grain sizedistribution; the grains appear in different grey level and areseparated by a dark contrast line that indicate the grain boundary orinterface location. The bar scale at the bottom of each imagecorresponds to 100 micrometers.

FIG. 2B is a phase contrast image at high magnification with the phasesindicated in different grey levels; pores (201) are black, aluminumtitanate (202) is intermediate grey, alumina (203) is dark grey,feldspar (204) is light grey.

FIG. 2C shows only the aluminum titanate grains with a separation of theindividual grains by an arbitrarily chosen grey level. A 1.5 degreec-axis mis-orientation was chosen as a distinction criterion betweendifferent grains.

FIG. 2D shows a larger area and represents only the aluminum titanatephase with a separation of individual domains by an arbitrarily chosengrey level. A 15 degree c-axis mis-orientation was chosen as distinctioncriterion between different domains.

The black scale bar at the bottom of the images corresponds to 500micrometers.

FIGS. 3A and 3B shows SEM images at different magnifications of polishedcross sections of fired ceramic lithium-free materials A. FIG. 3A showsthe pore structure (the bar scale is 200 micrometers). In FIG. 3B thephase distribution is shown with pores being black regions, aluminumtitanate being grey, alumina being dark grey, feldspar being light grey,and the microcracks being black wavy lines that clearly propagatethrough aluminum titanate grains and also through the feldspar.

FIGS. 3C and 3D show the material B having 0.5% lithium oxide in thebatch after firing at 1410 C. FIG. 3C shows the pore structure (the barscale is 200 micrometers). FIG. 3D shows the phase distribution withpore being black, aluminum titanate being mid-grey, alumina being darkgrey, feldspar being light grey with small glass-filled grain junctions,and the microcracks being black wavy lines that propagate preferentiallythrough the aluminum titanate grains and the intergranular glass film.

The distribution of pores in the materials with lithium is very similarto that of the lithium-free material, FIGS. A and C. Differences arevisible in the phase distribution in FIGS. 3B and 3D; the glass pocketsin the lithium-free material are much larger in size than in the lithiumcontaining material. Extension of the fine glass pockets in the lithiumcontaining materials along the grain boundaries can be seen. Microcracksin the lithium-free material are very large and cross without preferencefeldspar and aluminum titanate grains. In the lithium-containingmaterial, statistics show much finer cracks and also indicate that thecracks preferentially propagate through aluminum titanate grains andinterfaces and not through bulk feldspar grains.

XRD of the fully fired (1410° C., 15 h) materials with Li do not showany additional crystalline phases besides aluminum titanate,Sr-feldspar, unreacted excess alumina, and traces of rutile. The phasefractions are also similar to that in Li-free batches. Firing at toptemperature 1400° C. for 15 h resulted in the same phase composition,demonstrating the possibility of firing the materials with additives atlower temperatures. XRD (not provided here in a detailed listing) offully fired material C containing 1% Li after firing at top temperature1410° C. for 15 h and material B containing 0.5% Li after firing at toptemperature 1410° C. for 15 h/1410° C./15 h firing at top temperature1400° C. for 15 h (4 C), respectively, confirms that full conversioninto the final product was reached in all cases. XRD shows aluminumtitanate and strontium feldspar as main phases and only typical minorlevels of alumina and titania. Even at the lower firing temperature of1400° C., material C reaches the full conversion and phase compositionthat is achieved for the lithium-free material only after firing of atleast 15° C. higher temperature.

Physical Properties.

Physical properties of fired ware without any Li₂O addition (materialA), with 0.5% (material B), and 1% wt Li₂O (material C) are listed inTable 2.

TABLE 2 Fired fully fired 1410° C./ % d50 (d50 − CTE (RT to MOR 15 hrsporosity (microns) d10)/d50 1000° C.) (psi) Material A 49.5 15.4 0.524.3 250-280 without Li Material B 50.34 15.7 0.42 3 322 with 0.5% LiMaterial C 50.15 16.7 0.36 3.3 333 with 1% Li

Porosity.

The ware made from the Li-containing batches demonstrates advantages inits porosity compared to the Li-free reference batch, see FIG. 4. FIG. 4shows comparable pore size distributions for materials A, B, and Chaving, respectively, 0%, 0.5%, and 1% Li-addition. The insert shows theLi-free material A having a slightly enlarged (shoulder) small pore sizecontent compared to materials obtained from lithium-sourced materialbatches B and C. The overall porosity is very similar, but the mediumpore size increases with increasing Li-content and the d-factordecreases with increasing Li-content. Details of the pore sizedistribution, as shown in the inlet of FIG. 4, show that fine porositywith d<5 micrometers is almost completely suppressed in the batch with1% Li-addition and that porosity with d<10 micrometers is about half ofthat in a Li-free batch. The pore size distribution is narrower formaterial made from Li-containing batches, indicating a lower pressuredrop for the more homogeneous pore size.

Thermal Expansion.

The thermal expansion behavior of the B-material batch containing 0.5%Li is unmodified compared to the Li-free reference A-material. TheC-material containing 1% Li shows a slightly enhanced hysteresis uponthermal cycling. Details are shown in FIG. 5A. The CTE in the range fromroom temperature to 800° C. is lower for the batches containing Li, 3.3and 3×10⁻⁷ K⁻¹, respectively, compared to the Li-free reference batchwith 4.6×10⁻⁷ K⁻¹. The stability upon thermal cycling is a generalconcern for batches and needs to be addressed in the presence ofadditives that may form glasses that can recrystallize during furtherheat treatment(s). To test the stability of the material upon cycling,the material with 1% Li was extensively cycled and then the CTEremeasured, FIG. 5B. The hysteresis decreased upon cycling, the overallCTE increased slightly to 4.7×10⁻⁷ K⁻¹, indicating that glass pocketspartially crystallize during the heat treatment. A variation in CTE of1×10⁻⁷ K⁻¹ is within the typical variation of standard materials.

FIGS. 5A and 5B show thermal expansion curves for flux-free referencematerial A and materials B and C made from batches having 0.5% and 1% wtLi: initial cycle (5A) and after five (5) temperature cycles (5B),respectively. FIG. 5A shows thermal expansion curves during heating fromroom temperature to 1000° C. and subsequent cooling back to roomtemperature for Li-free reference material A (alternating short and longdashed lines) and material B (short dashed lines) and material C (solidlines) made from batches having 0.5% and 1% lithium oxide source,respectively. The set of upper lines having a higher CTE correspond tothe cooling curve. The lower lines correspond to the heating curve. FIG.5B compared the thermal expansion curves during heating and cooling formaterial C as-fired (continuous lines) and after thermal cycling(crossed hatched grey lines). The set of upper lines with higher CTEcorrespond to the cooling curve, and the lower lines to the heatingcurve.

Elastic Modulus.

The elastic modulus and its hysteresis in response to thermal cyclingshow differences compared to Li-free material. To ensure that thosedifferences are not due to slight difference in porosity, the measureddata are shown together with data normalized on cell geometry andporosity. Another standard material with very different porosity wasadded to the graph to show that this correction effectively corrects forcell geometry and porosity differences, since the two standard materialsin their extrapolation to dense material show no difference. TheLi-containing material, however, shows higher E-mod values and a largerhysteresis. Compared to other examples where simply higher microcrackdensity in the material leads to a larger hysteresis, a major differenceof the hysteresis of widened temperature is noted. While the top E-modvalues of standard and 0.5% Li containing materials are the same, thehysteresis for the 0.5% Li containing material is wider in temperature.The extension of the hysteresis at high temperature suggests thatcomplete microcrack healing requires higher temperatures than in thereference material. This can be explained by, for example, interactionsof microcracks with the glass phase, glass phase penetration into thecracks, and inhibition or slowing of closure. The 1% Li material showsthe same features, together with an increase in microcrack density.

FIGS. 6A and 6B show elastic modulus hysteresis upon thermal cyclingwhere the lower lines correspond to the heating curves, and the upperlines correspond to the cooling curves. FIG. 6A shows the elasticmodulus upon thermal cycling for the lithium-free reference material A(squares), material AA (diamonds) having different porosity levels, andmaterial C (dots) made from a batch having 1% Li. FIG. 6B shows the adensity-normalized elastic modulus (i.e., elastic modulus divided bymaterial relative density) upon thermal cycling for lithium-freereference material A (squares), material AA (diamonds) having differentporosity levels, and material C (dots) made from a batch having 1 wt %Li. The results indicate that the curves for the lithium free materialhas curves which overlap and are different from the lithium containingmaterials.

Material Strength.

Based on the above information that the number of microcracks appears tobe increased, but their length confined to the aluminum titanateagglomerates, and that Li-containing material has the same porosity asLi-free material and a slightly larger pore size, a direct comparison ofmodulus of rupture (MOR) as determined by flexure tests can be made. Thecomparison indicates that the material strength increases withincreasing Li-content from the standard Li-free material A with about260 psi to 320 psi for material B with 0.5% Li and to 330 psi formaterial C made with 1% Li. This is a significant increase in MOR ofmore than 25% for a minor modification in chemical composition.

Conclusion on Li-Oxide Sinter Additive.

Since all other physical properties of the material have been slightlyimproved in presence of small amounts of Li, the strength improvement atsuch very low additive levels and the suggestion of a lower firingtemperature than reference material A provide a significant advantagefor an improvement of known aluminum titanate-feldspar composites, sinceit provides a path to make higher porosity, thinner wall, dieselparticulate filters with significantly higher strength. Additionally oralternatively, the addition of lithium oxide permits a reduction ineither the length of the firing cycle or the top firing temperature.

Example 2

Phosphorous Oxide Glass-Former.

Since phosphorous oxide is known to promote glass formation in variousoxides, AT-type batches with phosphorous oxide levels from 1 to 5% weremade. Phosphorous oxide was added to the batch in form of aluminumphosphate. Batch compositions are summarized in the Table 3.

TABLE 3 Phosphorous oxide containing batch compositions Batch (wt %) A DE F INORGANICS Silica 10.19 10.19 10.19 10.19 Strontium Carbonate 8 8 88 Calcium Carbonate 1.38 1.38 1.38 1.38 Titanium Dioxide 29.95 29.9529.95 29.95 Hydrated Alumina 3.71 3.71 3.71 3.71 Lanthanum Oxide 0.2 0.20.2 0.2 Alumina 46.57 46 45.57 44.57 Aluminum Phosphate (purity 95%) 0 12.3 4.7 Totals 100 100.43 101.3 102.7 PORE FORMERS Graphite 10 10 10 10Potato Starch 8 8 8 8 SOLID BINDERS/ORGANICS Methylcellulose 4.5 4.5 4.54.5 OTHER LIQUID ADDITIONS Emulsion T 16 16 16 16

The fully fired alumina titanate materials with 1 to 4.7% AlPO₄,materials D to F, show the same feldspar, aluminum titanate, and aluminadistribution as the phosphorus-free material A, but are distinguishedfrom it by the presence of glass pockets. At high phosphate levels, theresulting microstructure is completely altered. The phosphorus isincorporated into the feldspar under formation of mixedsilicate-phosphate feldspar. Some aluminum titanate phase is dissolvedin the phosphorous-containing glass. Exchange reactions, as visible inthe SEM images of the materials by a change in contrast, have occurredat the borders of both the aluminum titanate grains and the feldspargrains. Microcracks preferentially follow the aluminum titanate graininterfaces and do cross, with less probability, the aluminum titanateand the feldspar bulk grains. As the phosphate content increases, boththe microcrack size, and the level of unreacted alumina increase.Samples having higher levels of phosphate show varying levels ofresidual glass depending on their cooling, post-annealing cycles, orboth. In samples with low levels of phosphorus, the low melting glassdistributes uniformly across the microstructure, penetrates the pores,and promotes the sintering and disappearance of small pores. Thisresults in a narrower pore size distribution and a larger medium poresize.

FIGS. 7A to 7C, respectively, show composited SEM micrographs aspolished cross-sections of three fired ceramic samples of batches D, E,and F, having three levels of magnification (left=low, middle=medium,right=high) and having an increasing amount of added aluminum phosphate,respectively: 1% (FIG. 7A), 3% (FIG. 7B), and 5% (FIG. 7C). While thelowest level of phosphorus addition, FIG. 7A, does not induce any majormodification in the material microstructure compared to P-free materialA and excels only by thin intergranular glass films and small glasspockets, the materials with higher phosphorus level in FIGS. 7B and 7Creflect strong modifications in their microstructure with roundedaluminum titanate grains that have partially dissolved in the glass anda reacted feldspar that has incorporated high levels of phosphorus andsurrounds the aluminum titanate grains like a liquid or glassy matrixmaterial that recrystallized upon cooling. Extended phosphorusinter-diffusion zones can be noticed at the grain boundaries andinterfaces, where phosphorus was more slowly incorporated into the bulkgrains. The low magnification visualizes the pore size distribution(left) and at higher magnification (middle and right) visualizes thephase distribution.

Differential scanning calorimetry (DSC) shows a melting event uponheating and recrystallization upon cooling which both reflect thepresence of the glass phase. The glass in presence also decreases thefinal melting of the composite from greater than 1440° C. to about 1380°C., as shown by the DSC of green ware containing AlPO₄. Melting of theglass phase occurs at 1320° C. for 4.7 wt. % AlPO₄ addition and isvisible as an additional endothermic event in the heating DSC; thecooling DSC from top firing temperature to room temperature shows asmall exotherm that is related to the crystallization of the glass. Thecrystallization is faster in presence of larger phosphorus levels.During its melting, the glass wets grain boundaries and penetrateseasily into small pores where it promotes sintering and loss of finesize porosity.

FIGS. 8A and 8B show the differential scanning calorimetry results forselected fired materials A, D, and E. FIG. 8A shows the heating curveswith glass melting events that are visible as exotherms above 1200° C.for curves of materials E and F and are not visible for the referencematerial A without phosphorus. FIG. 8B shows the cooling curves ofmaterials A, E and F with glass crystallization occurring in materials Eand F as an endothermic event upon cooling around 1100° C. The P-freematerial A does not show such an endothermic event upon cooling. FIG. 8Ashows the typical characteristics of the formation of a glass phase thatmelts and recrystallizes upon cooling.

XRD measurements of the fully fired (1410° C., 15 hr) samples with AlPO₄did not show any additional crystalline phases besides aluminumtitanate, alumina phosphate, Sr-feldspar, unreacted excess alumina, andtraces of rutile. The feldspar phase is modified from a triclinicstructure to a monoclinic structure. Glass levels are low and notvisible in the XRD.

Physical properties of fired ware A without any AlPO₄ addition andmaterials D, E, F with 1 to 4.7% AlPO₄ super-addition of phosphate arecompared in Table 4.

TABLE 4 Fired at 1410° C./ Porosity d₅₀ in d- CTERT-1000° C. in MOR in15 hrs. in % microns factor 10⁻⁷ K⁻¹ psi A 51.2 14.7 0.46 4.6 263 D - 1%46.8 18.6 0.28 8.5 352 AlPO₄ E - 2.3% 41.7 21.3 0.26 0.7 315 AlPO₄ F -4.7% 30.6 22.5 0.29 −1.2 384 AlPO₄

Ware made from batches containing AlPO₄ demonstrates a medium pore sizeincrease with increasing AlPO₄-content and decreasing width of the poresize distribution with increasing AlPO₄ content. Details of the poresize distribution show that the fine porosity with diameter less than 7micrometers is almost completely missing in batches with AlPO₄. Largermedium pore size and narrower pore size distribution promote a lowerpressure drop of the DPF.

FIG. 9A shows the pore size distribution properties of fired samples A,D, E, and F, made from batches having different levels ofphosphorous-source addition. The medium pore size has significantlyincreased with increasing phosphorus level (maximum shifted to largersize) compared to phosphorus-free material A. FIG. 9B shows an expandedview of the small pore region (arrow) of FIG. 9A illustrating a shoulderregion having preserved a higher small pore size content in material A(control, 0% aluminum phosphate) compared to the P-containing materialsD, E, and F that have lost their small pores through more effectivesintering in the presence of the phosphorus glass phase.

FIG. 10 shows thermal expansion curves for samples derived from batchesA, D, E, and F of Table 4 having different levels of phosphate addition.The thermal expansion behavior of the sample made from a batchcontaining 1% AlPO₄ is unchanged compared to the AlPO₄-free referencebatch. Samples with higher AlPO₄-levels have a smaller global expansionbetween room temperature and 800° C., but also have larger hysteresis.In addition they reflect a discontinuity in the curves that is due to aphase transformation of the mixed phospho-silicate-feldspar phase withtemperature. Although not bound by theory the larger microcracks arelikely responsible for the larger hysteresis.

As-measured MOR is presented in property Table 6. However, due tochanges in porosity and d₅₀, the porosity-normalized MOR provides moreinsight and shows a 23% increase in porosity-normalized strength foraddition of 1% AlPO₄, while higher levels of phosphate reverse thisbenefit due to the strong modifications of the feldspar and aluminumtitanate phases, and the large amount of glass. Porosity and poresize-normalized MOR (measured MOR normalized on material density andpore size) shows a doubling in strength.

It can be concluded that a small amount of alumina phosphate, such asabout 1% weight, can increase the strength of an AT-type material by asmuch as 34% and can have a negligible impact on CTE. Formation of smallamounts of phosphorous-containing glass causes only a small loss of theoverall porosity, eliminates the small-size pores, and decreases thewidth of the pore size distribution. The modified pore structure can beexpected to yield a lower pressure drop. The uniform distribution of thephosphorous-containing glass along grain boundaries and interfacespromotes crack propagation along the glass film and promotes crackdeflection, interface debonding, and crack multiplication within theglass, so that a stronger material is obtained (as demonstrated by thehigher MOR). However, the amount of glass has to be low. High levels ofaluminum phosphate addition yield formation of high quantities of glassphase and mixed feldspar and do not provide the thin grain boundaryglass layers as a preferred crack propagation path and result in lowermaterial strength.

Example 3

Glass formation by dipping fully fired honeycomb in phosphoric acid. A,G, and H fully fired parts were dipped into phosphoric acid (10% wt.aqueous solution) and then annealed (1400° C.). These batch compositionsare listed in the Table 5.

TABLE 5 Batch compositions of materials used for the phosphoric aciddipping. Batch (wt %) A G H INORGANICS Silica 10.19 8.75 8.75 StrontiumCarbonate 8 8 8 Calcium Carbonate 1.38 1.38 1.38 Titanium Dioxide 29.9529.95 29.95 Hydrated Alumina 3.71 3.71 3.71 Lanthanum Oxide 0.2 0.2 0.2Particulate alumina 46.57 — — Fibrous alumina — 58.01 48.01 Total 100100 100 PORE FORMERS Graphite 10 Potato Starch 8 20 20 SOLIDBINDERS/ORGANICS Methylcellulose 4.5 6.5 6.5 OTHER LIQUID ADDITIONSEmulsion T 16 20 20

The H₃PO₄-dipped and fired parts show a decrease in porosity with anelimination of fine pores and narrowing of the pore size distributioncompared to the original ware. The dipping results also in shrinkage ofthe parts, e.g., 8.5% in diameter and 5.3% in length. A strong decreasein porosity was observed for high porosity parts G and H, while lowerporosity A parts almost preserved their porosity. MOR increased in allcases, for G, H, and A by 51%, 97%, and 29%, respectively. Normalizationon porosity provided an increase of porosity-normalized MOR by 7%, 35%,and 28%, respectively.

Table 6 summarizes the properties of as-fired ware A, G, and H, and thesame fired ware that was dipped (“-d”) in phosphoric acid and thenannealed A-d, G-d, and H-d.

TABLE 6 CTE (RT to d50 in MOR/(1 − Treatment MOR 1000° C.) porositymicro- (d50 − porosity) Ware Conditions in psi in 10⁻⁷K⁻¹ in % metersd10)/d50 Permeability in psi G fired at 1410° C. 310 8.4 61.7 14 0.53532 810 G-d dipped in 467 7.6 46.1 17.4 0.28 702 867 10% H₃PO₄ H firedat 1410° C. 220 4.5 61.1 17.4 0.51 832 565 H-d dipped in 434 3.1 42.920.0 0.24 934 760 10% H₃PO₄ A fired at 1410° C. 216 3.4 48.3 15.8 0.41549 418 A-d dipped in 279 9.1 48.0 16.8 0.32 671 536 10% H₃PO₄

FIGS. 11A to 11C show pore size distributions for fired samples G (11A),H (11B), and A (11C), respectively, prior to dipping (diamonds) andafter dipping (squares) (“-d” suffix) the fired samples in H₃PO₄, andthen annealing.

FIG. 11A shows the pore size distribution of for as-fired material G andmaterial G-d that was obtained by dipping fired material G in 10% H₃PO₄and annealing it at 1400° C. The pore size distribution of this highporosity material after dipping and annealing is much narrower, allpores with diameter smaller than 10 micrometer in the mercuryporosimetry measurement have completely disappeared and the medium poresize has been increased.

FIG. 11B shows the pore size distribution for as-fired material H andmaterial H-d that was obtained by dipping fired material G in 10% H₃PO₄and annealing it at 1400° C. The pore size distribution of this highporosity material after dipping and annealing is much narrower, allpores with diameter smaller than 10 micrometer in the mercuryporosimetry measurement have completely disappeared and the medium poresize has been increased.

FIG. 11C shows the pore size distribution of for as-fired material A andmaterial A-d that was obtained by dipping fired material G in 10% H₃PO₄and annealing it at 1400° C. Even though the pore size distribution isstill narrower after dipping and annealing, the effect is lesspronounced for this lower porosity material than for materials G and H.

XRD results (not shown) of A-d, G-d, and H-d, i.e., after dipping thefired samples A, G, and H into 10% phosphoric acid and annealing them at1400° C., still show aluminum titanate and feldspar as main phases, butindicate a change in the symmetry of the feldspar phase from triclinicto monoclinic.

TABLE 7 XRD Calc. Extrusion Aluminum feldspar Aluminum alumina- Sampletitanate Alumina Rutile Feldspar type phosphate phosphate G 65.4 13.61.3 19.7 triclinic — G-d 66.4 15 1.1 15.9 monoclinic 1.7 5.7 H 74.3 4.20 21.1 triclinic — H-d 75.1 5.4 0 17.6 monoclinic 1.4 6.4 A 72.2 5.8 0.521.6 triclinic — A-d 74.6 7.4 0 15.5 monoclinic 1.9 3.8

FIGS. 12A to 12F show SEM images of polished cross sections (FIG. 12A to12C) and surfaces (FIG. 12D to 12F) of as-fired parts and these sameparts after being dipped in phosphoric acid and then annealing. FIG. 12Acorresponds to as-fired sample G (undipped; left) and G-d (dipped;right). FIG. 12B corresponds to as-fired sample H (undipped; left) andH-d (dipped; right). FIG. 12C corresponds to as-fired sample A(undipped; left) and A-d (dipped; right). FIG. 12D corresponds to thesurface of as-fired sample G (undipped; left) and G-d (dipped; right).FIG. 12E corresponds to the surface of as-fired sample H (undipped;left) and H-d (dipped; right). FIG. 12F corresponds to the surface ofthe as-fired sample A (undipped; left) and A-d (dipped; right). Allphosphoric acid-dipped samples show some modification of the feldsparphase compared to the starting materials that result from phosphorousglass that penetrates during annealing along the grain boundaries andinterfaces during the post-annealing, and dissolves the feldspar andaluminum-titanate grains while expanding the amount of glass. Part ofthe post-anneal glass phases crystallize upon cooling after thepost-annealing. The microcracks after dipping and annealing are alsomodified; they follow the AT grain boundaries and are deflected at thefeldspar/glass intersections.

While the average CTE is not changed much by the dipping procedure, theincreased microcrack density after dipping causes a much largerhysteresis in the heating-cooling cycle for dipped and post-fired partsthan in undipped parts.

FIGS. 13A, 13B, and 13C show CTE of original, phosphoric acid dipped,and post-fired parts of A, G, and H, respectively. The increase inmicrocrack density after phosphoric acid dip is also visible in theenhanced thermal hysteresis of the E-mod.

FIGS. 14A, 14B, and 14C, respectively, show the hysteresis of theelastic modulus during heating and cooling of fired-parts and phosphoricacid dipped fired-parts for samples A, H, and G.

Dipping of H parts into a 10% solution of phosphoric acid followed bypost-firing at 800° C. or 1000° C. resulted in no alteration inporosity, and an increase in MOR. During firing to 1400° C., a meltingevent occurred, accompanied by shrinkage and considerable loss ofporosity. Table 8 lists the porosity modifications.

TABLE 8 Modification in porosity for MTP as a function of post-annealingtemperature Post- Fully Acid annealing d50 (d50 − CTE (RT to firedExtrusion dip temperature % (micro- d10)/ 1000° C.) MOR Sample Condition(C.) porosity meters) d50 (in 10⁻⁷K⁻¹) (psi) G — 61.74 14.2 0.53 8.4 310G-d800° C. 10% 800 61.58 13.5 0.59 H₃PO₄ dip G- 10% 1000 59.30 13.8 0.57195 d1000° C. H₃PO₄ dip G- 10% 1400 46.13 17.4 0.28 7.6 467 d1400° C.H₃PO₄ dip

A silica-phosphate glass is formed during post-annealing after dipping,reacts with the feldspar, and dissolves alumina and titania at theperiphery of the aluminum titanate grains. The three glass compositionslisted in Table 9 were identified by microprobe analysis. The threeglass compositions contain high levels of silica with strontium,calcium, phosphorous, alumina, and dissolved titanium, which matches amodified feldspar with compositions of CaAl₂PSiO₈, SrAl₂PSiO₈, orNaAl₂PSiO₈. The glass promotes sintering, eliminates small pores, andnarrows the pore size distribution. Once again, the glass yields higherstrength; however, the MOR increase achieved by the dipping process issmaller than that obtained by addition of aluminum phosphate to thebatch.

TABLE 9 Phosphorous-rich glass compositions in weight % detected bymicroprobe in wares after dipping in H₃PO₄ and then annealing at 1400°C. SrO SiO₂ CaO TiO₂ P₂O₅ Al₂O₃ La₂O₃ 1 11.92 39.67 3.91 6.32 13.8920.29 0.49 2 11.11 39.57 4.60 6.55 17.34 19.05 0.58 3 10.02 40.63 3.903.85 15.52 23.79 1.85

Example 4

Materials I, J, K, and L have been batched to contain aluminum titanateand strontium feldspar with residual alumina with additional titania,strontia, lanthanum oxide, and silica to provide 0, 1, 2, and 5% oftitanium-strontium-lanthanum silicate glass. An additional batch M wasmade without any excess alumina and 5% of the same glass.

TABLE 10 Batch in wt % I J K L M INORGANICS Silica 8.14 8.48 9.17 9.8610.89 Strontium Carbonate 7.88 7.91 7.96 8.01 8.78 Calcium Carbonate1.38 1.41 1.48 1.54 1.85 Titanium Dioxide 31.72 31.53 31.17 30.81 31.50Lanthanum Oxide 0.03 0.08 0.13 0.13 Alumina 50.88 50.64 50.15 49.6646.87 Total 100 100 100 100 100 PORE FORMER Superaddition Potato Starch15 15 15 15 15 SOLID BINDERS/ORGANICS Methylcellulose 4.5 4.5 4.5 4.54.5 OTHER LIQUID ADDITIONS Emulsion T - 16 16 16 16 16

TABLE 11 CTE (RT Medium to 1000° C.) MOR Porosity pore size (d50 − Batch% Glass in 10−7K−1 in psi in % in microns d10)/d50 I 0 16.9 241 51.5217.83 0.44 J 1 10.6 211 52.11 17.83 0.44 K 3 6 238 50.48 17.21 0.42 L 55.4 270 49.89 18.23 0.38 M 5 3.3 252 45.62 21.45 0.31

While the addition of glass does not introduce any loss in porosity andin median pore size, it does impact the strength of the sample.As-measured MOR and normalized MOR show a significant increase withincreased glass level in the sample. The zero glass reference does notfollow the trend, because a “zero” glass level sample is difficult tobatch due to natural impurities in the raw materials and thereforeresulted in a slightly modified phase mixture.

FIG. 15 shows pore size distribution results for aluminumtitanate-strontium feldspar composite samples I, J, K, and L having 0,1, 2, and 5 wt % glass content. Despite significant changes in the glasslevel, only minor changes occur in the pore size distribution. Nodecrease in median pore size due to enhanced sintering in presence ofthe glass phase was observed. For 5% of glass, even a small increase inmedian pore size compared to the materials with lower or no glass levelcan be noticed. The 5% glass containing material in addition shows asignificant decrease in its fraction of small size pores compared to theother materials, thus promising less pressure drop due to a narrowerpore size distribution.

FIGS. 16A and 16B show bar charts of as-measured MOR (16A) andporosity-normalized MOR (16B) of selected aluminum titanate-strontiumfeldspar composites having excess alumina, 0, 1, 3, or 5% glass, and acomposite without excess alumina phase and 5% glass. It can be difficultto batch a material with zero glass content. Small impurities can bepresent and can induce formation of small levels of undesired secondphases, including non-defined glasses. Therefore, the material with azero glass level will be excluded from the comparison. It does not haveexactly the same phase distribution as the other materials. For theother materials with 1-5% glass content, a continuous increase in MORwith the glass content can be observed. The increase in MOR may beattributed to the toughening mechanisms that are introduced by the glassfilms and contribute to decrease the subcritical crack growth during thefour point flexure MOR test, so that the strength of the material isimproved. Imaging of the materials has shown that cracks stronglyinteract with the glass grain boundary film, choose the glass film aspreferential crack propagation path, and interact strongly at graininterceptions with branching glass at triple points.

Example 5

Addition of boron oxide for glass formation in AT. Table 12 lists lowlevels (e.g., 0.5 to 2 wt %) of boron oxide that were added as glassformer to a cordierite-mullite-aluminum titanate composite having 6 wt %MgO, batch N. Batch O contains 0.5%, batch P 1%, and batch Q 2% hasboron oxide.

TABLE 12 Batch Batch Ingredient Wt % N O P Q alumina 45.1 45.1 45.1 45.1silica 14.7 14.7 14.7 14.7 titania 34.1 34.1 34.1 34.1 magnesia 6.1 6.16.1 6.1 total inorganics 100% 100% 100% 100% superaddition B₂O₃ 0 0.5 12 superaddition pore 20 20 20 20 former

XRD (not shown) of the samples shows as a main phase (Mg, Al, Ti)₃O₅ andmullite; cordierite is present as minor phase. While the main phases donot change with addition of boron, the cordierite phase fractiondecreases with boron addition, due to incorporation of silica in theglass phase.

Table 13 summarizes intermediate firing properties and properties of thefully fired samples of CMAT with 6% MgO and 0 to 2% addition of boronoxide.

TABLE 13 CTE At At fully fired (RT (d50 − 600° C. 800° C. compression %% d50 to d10)/ MOR MOR strength Batch B₂O₃ porosity (microns) 1000 C.)d50 (psi) (psi) (psi) N 0 48.3 15.3 13.2 0.28 1244 O 0.5 48.8 14.8 11.30.28 1122 P 1 35.4 15.6 15 0.18 60 (308 64 (158 1713 Humidity Humidityaged) aged) Q 2 34.9 16.7 27.1 0.17 75 (95 1802 Humidity aged)

The resulting properties listed in Table 13 had findings similar topreceding examples. With addition of boron oxide above 0.5%, the overallporosity in fired materials decreased, while the median pore sizeincreased and the d-factor decreased as a result of promoted sinteringin the presence of the glass phase and loss of small porosity. Thematerial strength apparently increases substantially with boron oxideadditive of greater than 0.5%. For the lowest level of boron oxide of0.5%, the additive level was too low to produce a significant glasslevel with a homogeneous distribution and good wetting of theboundaries. Additionally, the intermediate firing strength was stronglyimproved in presence of boron oxide. While boron oxide-free (or 0.5% andbelow) batches have such low strength that the firing strength cannot beeasily measured, samples having 1 and 2% show acceptable firing strengthof 64 and 75 psi or after humidity aging prior to firing of 158 and 95psi, respectively.

FIG. 17 shows pore size distributions of alumina-cordierite-mullitecomposites having 6% MgO, and having 0% (sample N; control), 0.5%(sample O), 1 (sample P), or 2 wt % (sample Q) of boron oxide additionin the batch. FIG. 17 shows the global pore size distribution and theFIG. 17 insert shows an enlarged shoulder in the region of small sizes,especially for sample N (0% boron oxide addition; control). It can beseen that compared to the glass-free material, the addition of analumino-titanio-lanthano-silicate glass helps to shift the median poresize to larger values and at the same time to suppress the small pores.While an increase in median pore size is only visible for 3 and 5% ofglass, the loss of small pores (shoulder in the insert for the glassfree batch) is efficiently realized already for 1% glass addition.

Characterization Techniques

Microstructure Characterization.

Standard scanning electron microscopy (SEM) characterization wasconducted on honeycomb wall surfaces and polished honeycomb crosssections (cut perpendicular to the honeycomb channels). For theobservation of polished sections, the fired ware was infiltrated withepoxy, sliced and polished. The spatial distribution of porosity andphases present at a microscopic level was visualized on polished samplecross sections. The material porosity PSEM in %, average pore sized_(average), SEM were evaluated by SEM in combination with imageanalysis techniques.

Orientation Mapping (SEM/EBSD).

Electron backscattered diffraction (EBSD) on the SEM was used fororientation mapping of polished sample sections to derive grain size,relative orientation, and texture of the phases present with respect tothe honeycomb geometry. The samples were embedded in epoxy, carefullypolished and coated with a thin (10 Å) layer of iridium. All EBSDanalysis was completed on a Hitachi SU70 SEM, equipped with anOxford/HKL EBSD system. Data collection was completed for areas ofapproximately 1900 micrometers×700 micrometers at 23×10-9 Ampere beamcurrent, 20 kV accelerating potential, an interstep of 2 micrometers foroverall orientation information and 0.2 micrometers for grain sizedetermination. Phases used for identification of fully fired ATmaterials included aluminum titanate, corundum, rutile, and feldspar(Ca, Sr aluminum silicate). Pole figures were generated using HKL Mambosoftware with 5° data clustering.

Porosimetry.

Pore size distributions were explored by mercury intrusion porosimetryusing an Autopore IV 9500 porosimeter. This method uses the capillarylaw with non-wetting liquid and cylindrical pores. It is typicallyexpressed with the Washburn equation D=−(1/P) 4y cos θ, where D is thepore diameter, P the applied pressure, y the surface tension, and θ thecontact angle. The volume of mercury is directly proportional to thepressure. Micrometrics software data reduction uses the differential andlog differential to calculate the first derivative of the cumulativespecific intrusion volume as a function of calculated log diameter.

Thermal Expansion.

Thermal expansion was measured for bar-shaped samples with dimension0.25″×0.25″×2″ during heating from room temperature to 1200° C. at arate of 40K/min and subsequent cooling to room temperature. For the datareported in the property Table 13, the long axis of the test bars wasoriented in the direction of the honeycomb channels, to provide thethermal expansion in the axial direction of the honeycomb parts.

Average thermal expansion coefficients for various temperature rangeswere recorded in the property Table 13, CTE₂₀₋₈₀₀ in K−1, the averagethermal expansion coefficient from room temperature to 800° C., definedas L(800° C.)−L(20° C.)/780 as average thermal expansion coefficient inthe temperature range from room temperature to 800° C., CTE₂₀₋₁₀₀₀ inK−1, the average thermal expansion coefficient from room temperature to1000° C., defined as L(1000° C.)−L(20° C.)/980 as average thermalexpansion coefficient in the temperature range from room temperature to1000° C., CTE₅₀₀₋₉₀₀ in K−1, the average thermal expansion coefficientfrom 500 to 900° C., defined as L(900° C.)−L(500° C.)/400 as averagethermal expansion coefficient in the temperature range from 500° C. to800° C. CTE₅₀₀₋₉₀₀ is of particular importance for the application ofhoneycomb parts for exhaust after treatments in the automotive vehicle,where the honeycomb part is subjected to severe rapid temperaturechanges and where the temperature range from 500 to 900° C. would matcha frequently encountered operation temperature range.

Mechanical Strength.

A ceramic's strength can be tested using three or four bending. Themaximum stress prior to failure is often referred to as the modulus ofrupture or MOR. Strength values (4-point flex, MOR) have been measuredusing four point flexure with a lower span of 2″ (50.8 mm) and an upperspan of 0.75″ (19 mm). The specimen geometry for the 4-point flexuretests was 2.5″ (63.5 mm) in length, 0.5″ (12.7 mm) in width and 0.25″(6.4 mm) thick. The force-measuring system used was equipped with aread-out of the maximum force and a calibrated load cell. The MOR valuewas calculated using the flexure strength equation.

$\sigma_{{4 - {point}},{MOR}} = {\frac{3}{4}\frac{PL}{{bd}^{2}}}$

However, this equation does not account for the cellular channelsthrough the specimen and is not the true strength of the material. Allspecimens tested had a square cellular (honeycomb) with the channels inthe direction of the length. The actual material strength, oftenreferred to as the wall strength (σ^(wall)), has to be determined byaccounting for the cellular structure.

Elastic Modulus Measurement.

Bar-shaped samples with dimension 5″×1″×0.5″ and the long axis beingoriented in the direction of the honeycomb channels were used to measurethe elastic modulus by flexural resonance frequency. Samples were heatedto 1200° C. and cooled back to room temperature. For each temperaturethe elastic modulus was directly derived from the resonance frequencyand normalized for sample geometry and weight by referring to ASTM C1198-01.

Strain Tolerance.

A strain tolerance MOR/E-mod can be derived from the strength of thematerial and its Young's modulus. This strain tolerance describes theability of a material to handle strain. The higher the strain tolerance,the less likely the material will fracture. The strain tolerance isindependent of the honeycomb geometry and can directly be compared forparts with different cell densities or wall thickness.

Thermal Shock.

The resistance to thermal shock for a honeycomb for use as an automotiveexhaust after-treatment is significant because the part undergoes severethermal cycling during rapid heat up, cool down, and duringregeneration. Thermal shock resistance of materials is often predictedby a figure-of-merit. The most common of these is the R parameter, whichis simply the temperature difference which will result in a stress thatexceeds the strength of the material. A second thermal shockfigure-of-merit is the R_(st) parameter that takes into account thepre-cracked state of a material, and is similar to the R parameterexcept that it takes into account the toughness (K) of the materialinstead of the strength. The thermal shock resistance of a honeycomb isexpected to improve with high strength, high toughness, low elasticmodulus, and low thermal expansion of its material.

$R = {{\frac{\sigma}{\alpha \cdot E}\mspace{76mu} R_{st}} = \frac{K}{\alpha \cdot E}}$

Any aspect, feature, or embodiment of the foregoing disclosure can beused in any combination or permutation with any one or more otheraspect, feature, or embodiment recited in the appended claims.

The disclosure has been described with reference to various specificembodiments and techniques. However, many variations and modificationsare possible while remaining within the scope of the disclosure.

1.-11. (canceled)
 12. A method for improving the thermo-mechanical properties of an aluminum titanate-based composite, comprising: dipping a fully fired aluminum titanate composite into an aqueous phosphoric acid solution of 0.5 to 10 wt %; and annealing the dipped composite, to provide phosphorous incorporation into the resulting composite of 0.5 to 2 wt %, the weight % being based on a superaddition relative to the weight of the un-dipped composite.
 13. A method for toughening a microcracked aluminum titanate ceramic, comprising: creating an intergranular glass film within the ceramic, the film having a thickness of from about 20 nm to 500 nm, and the film interacts with the microcracks and limits uncontrolled growth of the microcracks.
 14. The method of claim 13 wherein the ceramic is toughened from about 5 to about 25%, as demonstrated by an increase in the modulus of rupture measured by 4-point bending relative to a ceramic prepared without the intergranular glass film.
 15. The method of claim 12 wherein the dipped and annealed aluminum titanate-based composite has an average CTE that is substantially unchanged by the dipping procedure, the microcrack density of the dipped and annealed aluminum titanate-based composite is increased, and the dipped and annealed aluminum titanate-based composite has a significantly larger hysteresis in a heating-cooling cycle compared to an undipped and annealed composite.
 16. The method of claim 12 wherein the dipping is accomplished in aqueous 10 wt % phosphoric acid.
 17. The method of claim 12 wherein the annealing is accomplished at 1400° C.
 18. The method of claim 12 wherein the aluminum titanate-based composite comprises grains of an aluminum-titanate crystalline phase and grains of a strontium-feldspar crystalline phase, and an intergranular glass phase between the grains of the aluminum-titanate crystalline phase and the strontium-feldspar crystalline phase.
 19. The method of claim 12 wherein the thermo-mechanical property of the dipped and annealed composite is material strength and the material strength is improved from 310 psi to 467 psi compared to an un-dipped composite.
 20. The method of claim 12 wherein the thermo-mechanical property of the dipped composite is the CTE which is decreased from 8.4 10⁻⁷K⁻¹ to 7.6 10⁻⁷K⁻¹ over 25 to 1000° C. for a dipped and annealed composite having a pore size increased by 3 micrometers compared to an un-dipped and annealed composite.
 21. The method of claim 12 wherein the resulting annealed aluminum titanate-based composite has a porosity greater than 42%, a median pore size greater than 16 micrometers, and a CTE of less than 9×10⁻⁷K⁻¹ from 25 to 1000° C.
 22. The method of claim 12 wherein the resulting annealed aluminum titanate-based composite is a honeycomb filter.
 23. The method of claim 12 further comprising selectively plugging the ends of the honeycomb filter in an alternating checkerboard pattern to form a through wall filter for exhaust gas particle filtration.
 24. The method of claim 22 further comprising washcoating the honeycomb filter with a catalyst suitable for catalytic conversion applications.
 25. The method of claim 12 wherein the feldspar phase, prior to annealing, has a monoclinic structure.
 26. The method of claim 12 wherein the annealing produces a strengthened product having a feldspar phase with a triclinic symmetry. 